Shaped articles comprising polycrystalline SiC are known. They are characterized by excellent physical properties, such as high resistance to thermal shock, abrasion and oxidation, together with high levels of strength and thermal conductivity. It is this combination of properties which makes SiC materials leading candidates for engineering applications. However, this combination of properties only concurs in high density materials.
During high temperature heat treatments of prerequisite powder compacts, a reduction in the surface energy of the system can occur. The reduction in surface energy is through the diffusion of atoms by either grain boundary diffusion and subsequent densification, or by grain growth through surface diffusion mechanisms with virtually no macroscopic densification. At the high temperatures required for the sintering of SiC powder compacts, surface diffusion typically prevails over grain boundary diffusion. This results in coarsening of the SiC grains in a powder compact with little macroscopic densification taking place.
The oldest process for production of dense articles of SiC is that of reaction sintering, in which silicon liquid or vapour is infiltrated into a compacted body of SiC powder and C. The Si reacts with the C to form SiC in situ which bonds the powder particles. However, this process typically leaves from 8 to 12 volume percent of free Si, which sets a maximum operating temperature of about 1300.degree. to 1400.degree. C. for the resultant densified article.
In more recent times, attention has been directed to the use of certain additives which promote grain boundary diffusion over surface diffusion for pressureless sintering of SiC. However, apart from B or certain compounds thereof, found to be effective in increasing grain boundary diffusion, there does not appear to have been any successful proposal, at least in terms of commercial utility. Moreover, even with use of B or a B compound, problems still exist.
In use of B or a B compound, C usually is added as disclosed in U.S. Pat. Nos. 4,004,934, 4,041,117 and 4,108,929 all to Prochazka and U.S. Pat. No. 4,124,667 to Coppola et al. It is indicated that the C reduces the surface SiO.sub.2 layer on the SiC powder to SiC and CO. In U.S. Pat. No. 4,041,117, Prochazka suggests that the SiO.sub.2 can halt densification of SiC compacts so that little or no shrinkage can occur. Prochazka also suggests that the addition of C can limit exaggerated grain growth during densification. However, he further indicates that grain growth can only be inhibited by strict control of temperature and pressure within narrow limits. Also, the final product usually contains C particles in the microstructure which can lead to degradation of mechanical properties of the product.
The literature on effective sintering aids for SiC powder, other than B or B compounds, is credited as having commenced with Alliegro et al, J. Amer. Ceram. Soc. 39 [11] 386-89 (1056). This reference discloses that 1% Al addition to .alpha.- or .beta.-SiC powder enables densification by hot-pressing to about 98% of the theoretical density. The .beta.-SiC powder was synthesised from a Si/C mixture, in which case, the Al usually was added to that mixture as oxide that was reduced during the synthesis. With use of .alpha.-SiC powder, the Al evidently was added as the metal powder. Alliegro et al report that Fe, Li, Ca and Cr also aided densification, but that Mg, Ta, Co, Ba, Mo, W, Sr and Cu were not beneficial whether used alone or with Al.
Artemova et al, in Neoroanicheskie Materialy, Vol. 10, No. 12, pp 2228-9, Dec. 1974, report on the preparation of a densified product by shock compression employing an explosive charge. Powdered SiC/Al.sub.2 O.sub.3 mixtures ranging from 10/90 to 90/10 mole percent, in 10 mole percent increments, were used and attained in excess of 98% of the theoretical density for the mixtures. This mode of densification, having some similarity to hot-pressing, suggests the suitability of Al.sub.2 O.sub.3 as an additive in SiC densification by more conventional procedures. However, Artemova et al report that it has not yet proved possible to densify mixtures of SiC/Al.sub.2 O.sub.3 at all by the usual methods.
Lange, J. Mater. Sci. 10 [1975] 314-320, reports on the hot pressing of both .alpha.- and .beta.-SiC powder with use of Al.sub.2 O.sub.3 as a densification aid. While only quite small additions of Al.sub.2 O.sub.3 were used, ranging from 0.01 to 0.15 volume fractions, densities up to and greater than 99% of the theoretical density were achieved. Densification was attributed to a liquid phase which formed at high temperatures. However, the use of Al.sub.2 O.sub.3, even at such low levels, was reported to result in large second phase streaks of Al.sub.2 O.sub.3 of up to several millimeters long and usually 10 to 15 .mu.m wide. For brittle materials such as ceramics, the presence of a flaw, such as a crack, pore or inclusion can result in stress concentration leading to failure. Streaks as reported by Lange would inevitably be deterimental to the physical properties of the densified SiC, as they greatly increase the defect size in the material.
It was speculated by Lange that the streaks of Al.sub.2 O.sub.3 were the result of laminar void spaces present in the cold pressed specimens. Possible solutions to eliminate or reduce the occurrence of the streaks was to employ a "sandwich" approach using layers of thinner bodies which, when compacted, formed thicker bodies. This technique would be limited to procedures such as hot pressing. Another technique proposed was grain growth of the SiC grains. Under industrial conditions, the presence of such voids is always possible with the probability of their occurrence increasing with increasing thickness of the component. Grain growth of SiC to aid the removal of such streaks may prove difficult to control in practice. Excessive grain growth is a problem associated with some of the techniques employed for pressureless sintering of SiC. This is considered to be a disadvantage in the use of Al.sub.2 O.sub.3 as a densification aid. No indication was given by Lange as to whether Al.sub.2 O.sub.3 would be an effective aid for the pressureless sintering of SiC.
Omori et al, J. Am. Ceram. Soc. 65 [1982] C-92, disclose the use of oxide additives, viz. Al.sub.2 O.sub.3 and Y.sub.2 O.sub.3, in the Pressureless sintering of .beta.-SiC powder. The oxides were used separately at 10 wt %, and in combination to a total of 10 wt % at ratios of 4:1, 3:2, 1:1, 2:3 and 1:4. Densification was achieved with 10% Al.sub.2 O.sub.3, but only with 4% shrinkage and a relative apparent density of about 75%. With decreasing Al.sub.2 O.sub.3 content, densification was enhanced to about 97% of the theoretical density at an oxide ratio of 1:1, but the level of the fired bulk density thereafter decreased and, with 10 wt % of Y.sub.2 O.sub.3 alone, no benefit was obtained over .beta.-SiC alone. Omori et al reasonably conclude that Al.sub.2 O.sub.3 enhances pressureless sintering despite its partial loss attributed to sublimation, but that Y.sub.2 O.sub.3 does not. However, the results do suggest that, to a degree, Y.sub.2 O.sub.3 improves the beneficial effect of Al.sub.2 O.sub.3. Omori et al report the loss of Al.sub.2 O.sub.3 on sintering, with a residue of this oxide being determined by chemical analysis but not by X-ray diffraction.
A more recent study by Negita, J. Am. Ceram. Soc. 69 [12] C-308-C-310 (1986), reports on the selection of suitable densification aids for the sintering of SiC. Using thermodynamic arguments, Negita reported that metal additives such as B, Al, Fe, Ni and Mg should be effective sintering aids for SiC and that this had been found to be the case experimentally. In relation to B, Al and Fe, this accords with the work of others, as detailed above. On the basis of the same arguments, Negita reports that metal oxides, including Al.sub.2 O.sub.3, BeO, Y.sub.2 O.sub.3 HfO and rare earth oxides, should be effective densification aids, and that this was borne out experimentally at least for Al.sub.2 O.sub.3, BeO, Y.sub.2 O.sub.3, La.sub.2 O.sub.3, Ce.sub.2 O.sub.3 and ThO.sub.2.
In contrast to the oxides listed in the previous paragraph, Negita reports that metal oxides including CaO, MgO and ZrO.sub.2 are indicated not suitable as they tend to decompose SiC. In addition, Negita suggests that the use of C with metal oxides is indicated as beneficial in the case of Al.sub.2 O.sub.3, BeO, Y.sub.2 O.sub.3, CaO, ZrO.sub.2, HfO.sub.2 and rare earth oxides.
The use of Al.sub.2 O.sub.3 as a densification aid in the pressureless sintering of SiC powder is disclosed in U.S. Pat. No. 4,354,991 to Suzuki et al. The proposal of this reference is to mould a mixture of an oxygen-containing Al-compound, which can be converted into Al oxide by heating in a non-oxidative atmosphere at a ratio of 0.5 to 35 wt % Al.sub.2 O.sub.3, with the remaining ceramic material substantially being SiC. Such moulded mixture is subjected to pressureless sintering in a non-oxidative atmosphere at 1900.degree. C. to 2300.degree. C. Despite the requirement that the oxygen-containing compound is one which can be converted into Al oxide, it evidently is envisaged that the compound can be Al oxide. However, a number of disadvantages, of which some are confirmed by our work on the pressureless sintering of mixtures of SiC and Al.sub.2 O.sub.3, are apparent from U.S. Pat. No. 4,354,991.
The fired bulk densities obtained by the teaching of U.S. Pat. No. 4,354,991 are relatively low, and also subject to substantial variation with firing conditions. Also, the sintering times are relatively long, ranging from a preferred minimum of 2 hours up to 24 hours, with 3 to 5 hours being typical even with relatively small samples. In a continuous process for densification of SiC powder, such reaction times would result in lower production rates. Furthermore, another problem exists in the preferment for control and maintenance of Al species in the firing furnace atmosphere for long periods of time required for sintering.
No mention is made in U.S. Pat. No. 4,354,991 of the formation of streaks of Al.sub.2 O.sub.3 as reported in the above-mentioned article by Lange, even though such defects are likely to be a characteristic of the use of Al.sub.2 O.sub.3 alone. As suggested by Lange, long soak times may be required to eliminate such streaks, and this possibly explains the relative long sintering times taught by U.S. Pat. No. 4,354,991. However, as indicated herein, the use of such sintering times is disadvantageous.
We have found that a further apparent characteristic of the use of Al.sub.2 O.sub.3 alone as a sintering aid for SiC powder is the tendency for zoning, particularly in the production of relatively large articles. That is, we have found that use of Al.sub.2 O.sub.3 alone has a pronounced tendency to produce a well densified outer layer enclosing an internal core which can exhibit substantially less densification. Where zoning occurs, the article is at least less than optimum. Also, internal stress due to the zoning can result in the article exhibiting cracks or, in extreme cases, the article can fail completely with the outer layer spalling.
The tendency for zoning with the use of Al.sub.2 O.sub.3 alone as a sintering aid for SiC powder, as taught by U.S. Pat. No. 4,354,911, is believed to be due to the difficulty of producing a sufficient volume percent of liquid phase at an appropriate temperature. This difficulty may also explain the tendency for streak formation as reported by Lange, or streak formation may exacerbate the difficulty in achieving a sufficient volume of liquid phase. As is known, efficient liquid phase sintering requires not only the formation of a liquid phase at a suitable temperature, but also the presence of that phase in a sufficient volume over a suitable temperature range.
In the proposal of U.S. Pat. No. 4,354,911, formation of a suitable liquid phase is not possible simply by melting of Al.sub.2 O.sub.3, except at extremely high temperatures. The melting point of Al.sub.2 O.sub.3 is about 2015.degree. C., while decomposition and loss by volatilization of decomposition products thereof commences below that temperature, as recognised by Suzuki et al and also taught by the above-mentioned article by Omori et al. Despite the sole addition of Al.sub.2 O.sub.3 as a sintering aid, SiO.sub.2 also is present as an impurity layer up to about 2 wt % on finely divided SiC powder (unless previously removed), and the SiO.sub.2 can facilitate the formation of a liquid phase at a temperature below the melting point of Al.sub.2 O.sub.3.
Reference to the phase diagram for the SiO.sub.2 -Al.sub.2 O.sub.3 binary system shows a eutectic composition at about 93% SiO.sub.2 --7% Al.sub.2 O.sub.3 which has a melting point at about 1595.degree. C. Thus, assuming that the rate of heating to the sintering temperature range of 1900.degree. to 2300.degree. C. is not excessive, solid-solid diffusion between the separate Al.sub.2 O.sub.3 and SiO.sub.2 can give rise to an initial small volume of liquid at temperatures above 1595.degree. C. Also, SiO.sub.2 melts at about 1730.degree. C. and, assuming that the SiO.sub.2 is not previously volatilized or decomposed, as tends to occur, a small volume of SiO.sub.2 -containing liquid phase can be formed above 1730.degree. C. and this can increase in volume by taking up Al.sub.2 O.sub.3 by liquid-solid diffusion.
In the method taught by Suzuki et al in U.S. Pat. No. 4,354,911, the lower level of Al.sub.2 O.sub.3 addition is 0.5 wt %, corresponding to an SiO.sub.2 to Al.sub.2 O.sub.3 ratio on the Al.sub.2 O.sub.3 rich side of the eutectic of the SiO.sub.2 -Al.sub.2 O.sub.3 binary system. That is, when allowance is made for 2.0 wt % SiO.sub.2 being high and 0.5 wt % Al.sub.2 O.sub.3 being a minimum, it is apparent that a best possible ratio is about 80% SiO.sub.2 : 20% Al.sub.2 O.sub.3. A lower SiO.sub.2 content or a higher Al.sub.2 O.sub.3 content rapidly advances that ratio away from the eutectic composition to increasily richer Al.sub.2 O.sub.3 contents. At only 2.0 wt % Al.sub.2 O.sub.3, the ratio is at least at the mid-point of the SiO.sub.2 -Al.sub.2 O.sub.3 phase diagram. At 4.0 wt % Al.sub.2 O.sub.3, the ratio is such that little, if any, liquid previously formed will remain, with further liquid then not being formed until a temperature of about 1840.degree. C. is achieved. That is, with an Al.sub.2 O.sub.3 content of at least 4.0 wt % Al.sub.2 O.sub.3, any liquid initially formed will be substantially lost, due to precipitation of a corundum or mullite solid phase having a melting point of about 1840.degree. C. However, given that an Al.sub.2 O.sub.3 addition of only 0.5 wt % still is on the Al.sub.2 O.sub.3 -rich side of the eutectic composition, at least a proportion of any initially formed liquid with less than 4.0% Al.sub.2 O.sub.3 additions will similarly be lost due to precipitation of corundum or mullite. These problems are further exacerbated by the tendency for SiO.sub.2 and Al.sub.2 O.sub.3 to decompose and to be lost by volitization of their decomposition products at temperatures approaching 1840.degree. C., making it very difficult to produce, or produce and retain, a significant volume of a liquid phase. Also, Al.sub.2 O.sub.3 present at a level significantly in excess of 4.0 wt % will not be able to be taken fully into solution below at least about 1840.degree. C., with the temperature at which this is possible rapidly increasing with the level of Al.sub.2 O.sub.3 addition to about 2015.degree. C. Moreover, if there is only alumina present, a liquid phase cannot be formed below the melting point of Al.sub.2 O.sub.3, that is, below about 2015.degree. C., and even then, a liquid will only form if some Al.sub.2 O.sub.3 is retained until that temperature is attained.
The precipitation of corundum or mullite from initially formed liquid may explain the streaks of Al.sub.2 O.sub.3 reported in the above-mentioned article by Lange. The streaks are referred to by Lange as suggesting a "frozen liquid". This may well have resulted from corundum or mullite precipitated from an initially formed liquid, and only partially remelted on heating at about 1840.degree. C.
For the temperature range of 1900.degree. C. to 2300.degree. C. taught by Suzuki et al in U.S. Pat. No. 4,354,911, and the addition of Al.sub.2 O.sub.3 alone at from 0.5 to 35% as a sintering aid for SiC, it therefore is extremely difficult to achieve a liquid phase at all, let alone one in a sufficient volume for efficient liquid phase densification. As the Al.sub.2 O.sub.3 level increases above 0.5%, the temperature at which fully liquid SiO.sub.2 and Al.sub.2 O.sub.3 is present also increases, and the volume of liquid able to be produced below the 1840.degree. C. solidus decreases. Particularly above about 4% Al.sub.2 O.sub.3, it can be necessary to use a temperature substantially above 1900.degree. C. in order to achieve any significant volume of liquid at all.
A further disadvantage of the proposal of U.S. Pat. No. 4,354,991 arises from the strong preferment for use of .beta.-SiC powder, rather than, .alpha.-SiC powder. .beta.-SiC is not as readily available as .alpha.-SiC as produced by the conventional Acheson process for the manufacture of SiC grit. That process accounts for a major portion of world-wide production of SiC and .alpha.-SiC is readily available and is a commodity traded on the world market.
In International patent specification PCT/US88/00040 (W088/05032), Fuentes discloses the pressureless sintering of SiC powder, using as a sintering aid a combination of Al.sub.2 O.sub.3 and CaO. Fuentes recognises that with use of Al.sub.2 O.sub.3 alone as a sintering aid for SiC, the liquid phase necessary for sintering is deficient in volume and/or forms too slowly. He therefore teaches use of a sintering aid mixture which produces a liquid phase at from 1815.degree. to 1855.degree. C. and comprises Al.sub.4 O.sub.4 C and Al.sub.2 OC. However this liquid phase, which also can be generated by use of Al.sub.4 O.sub.4 C and Al.sub.2 OC ab initio, itself forms at an excessively high temperature for optimum densification. In addition, as reported by Foster et al, J. Am. Ceram. Soc. 39 [1956] 1-11, Al.sub.4 O.sub.4 C, Al.sub.2 OC and Al.sub.4 C.sub.3 are very unstable towards both moisture and oxygen. The presence of these species in the product resulting from the process taught by Fuentes is very undesirable, and to be expected to greatly degrade the performance and severely limit the utility of the product. The process and product as disclosed by Fuentes therefore presents significant disadvantages.
In contrast to the prior art discussed above relating to the densification of SiC to produce bodies of high density approaching the theoretical density, the use of oxides for the bonding of SiC grits to form refractory bodies also has been considered. Thus, in U.S. Pat. No. 2,040,236 to Benner et al, the use of a bonding material of Al.sub.2 O.sub.3 together with either CaO, MgO or a mixture of CaO and MgO was considered for use in bonding SiC grit in producing a refractory body. Benner et al teach the heating in a non-oxidizing atmosphere of a pressed mixture of SiC grit and such bonding material. The heating was to a relatively high temperature, at which the bonding material softened to undergo incipient fusion. However, the rate of heating to temperature was rapid, such as about 35 minutes. Also, it is emphasised that the time at temperature was to be short so that, while sufficient to soften the bonding material, recrystallization of SiC could be avoided. Furthermore, the time at temperature was to be short so that the bonding material did not either decompose or react with the SiC.
The suitable SiC grit proposed by Benner et al ranged from 14 mesh to less than 80 mesh, but with coarse, medium and fine size fractions. Thus 40% was--40 mesh+36 mesh (ranging from less than about 1170 .mu.m to about 410 .mu.m); 10% was--40+70 mesh (ranging from less than about 370 .mu.m to about 190 .mu.m); and 50% was of--80 mesh (ranging down from about 180 .mu.m). While only the small sub-micron portion of the fine size fraction would be appropriate for densification as required by the prior art discussed above, Benner et al report production of a useful refractory compared with use in a similar context of other bonding materials. Their refractory is said to have been very dense and of lower permeability in that context. Microscopic examination (as applicable in 1932) is stated to have shown the product to exhibit pores only partially filled with bond material, while the refractory was permeable to gases. In this regard, the disclosure of Benner et al is devoid of any indication that macroscopic densification of the body occured. Also, the SiC particles of the grit, as confirmed by reference to it as a filler, in essence was bonded in a matrix of the bonding material, with the latter evidently remaining in essentially the proportion of, for example, 5 to 10% in which it was added to the mixture.
The teaching of Benner et al detailed in the preceding two paragraphs is appropriate for the bonding of SiC grit, but does not provide guidance relevant to liquid phase densification SiC powders. That is, they are seeking to produce refractories by bonding SiC grit particles in a matrix. The matrix acts in effect as a cement or glue (in the general sense of these terms) which encapsulates and isolates the SiC grit particles without decomposing or reacting with the SiC of these particles. In contrast, liquid phase sintering necessitates finer SiC powder of a compact being densified being taken into solution and subsequently precipitated, such as onto larger SiC grains, with the end product having clearly defined grain boundaries between SiC grains and any second phase. In effect, Benner et al teach use of a passive bonding material which softens to form a matrix, whereas liquid phase sintering requires the presence of an active liquid phase which is formed by the assistance of sintering aids.
The non-oxidizing atmosphere proposed by Benner et al was required to inert to both the SiC and the bond material. Carbon monoxide is indicated as being satisfactory relative to Al.sub.2 O.sub.3, MgO and CaO and their mixtures. However, where SiO.sub.2 was a principal constituent of the bond material, a more inert atmosphere such as nitrogen or helium was preferred.
Further, in U.S. Pat. No. 4,829,027 Cutler et al disclose liquid phase sintering of SiC with use of a rare earth oxide and Al.sub.2 O.sub.3 ; the rare earth oxide principally exemplified being Y.sub.2 O.sub.3 as in the Omori et al reference considered above. The disclosure of this reference emphasises the importance of attaining a liquid phase at a relatively low temperature, in achieving densification by pressureless liquid phase sintering of SiC, substantiating our findings in relation to a dissimilar system based on use of Al.sub.2 O.sub.3.
Finally, Japanese patent application 01230472, public disclosure No. 89-230472, by Kurosaki Refractories Co. Ltd., proposes the production of SiC sintered products using alumina/magnesia spinel (i.e. MgAl.sub.2 O.sub.4) as a sintering aid. Kurosaki teaches that when spinel alone is used as a sintering aid, magnesia will evaporate preferentially from the surface of the spinel powder grains, leaving grain surfaces covered with a layer of Al.sub.2 O.sub.3. During sintering, a liquid phase is said to form at temperatures of 1900.degree. C. and above; this being seen as beneficial in resulting in little likelihood of deterioration of the excellent high temperature characteristics inherent in SiC. In this regard, the teaching of Kurosaki is to avoid a liquid phase being formed at fairly low temperatures, a matter on which they are at variance with the clear teaching of Fuentes, Cutler et al and our research.
A disadvantage of the teaching of the Kurosaki proposal is the reliance on relatively expensive spinel as the sintering aid, particularly as in excess of 5 wt % spinel is necessary for optimum results. In this matter, the same disadvantage exists with the proposal of Cutler et al in their reliance on expensive rare earth oxides. However, further major disadvantages exist with the proposal of Kurosaki. The first is that arising from the loss of MgO to which they refer since, with increasing level of spinel, the resultant weight loss will be increased; with a possible maximum of about 9.9 wt % due to this factor alone at 35 wt % spinel. However, as made clear by the work of others considered above, and also substantiated by our findings, these weight losses are likely to be exacerbated by additional loss of SiO.sub.2. Al.sub.2 O.sub.3 and SiC. A further important disadvantage is that, due to the spinel grains becoming coated with Al.sub.2 O.sub.3, any liquid phase initially tending to form will require slow solid-solid diffusion, followed by dissolution of Al.sub.2 O.sub.3 and spinel, with this occurring to any significant extent in a reasonable time only at temperatures substantially above 1900.degree. C. This will lead to essentially the same problems in achieving a sufficient volume of liquid phase necessary for efficient liquid phase densification, as discussed above in relation to the teaching of Suzuki et al.